High Strength Cold Rolled Steel Sheet Having Excellent Shape Freezability, and Method for Manufacturing the Same

ABSTRACT

A high-strength isotropic cold-rolled steel sheet having excellent shape-fixability, and a method for manufacturing the same are disclosed. The steel sheet is made using a low-carbon steel comprising a little amounts of Ti to have an r of 1.3 or less, r approaching 1, and a low in-plane anisotropy index (Dr) of 0.15 or less, ensuring excellent shape-fixability so as to be suitable for automotive outer panels which are essentially subjected to deformation in a stretching mode.

TECHNICAL FIELD

The present invention relates to a high-strength cold-rolled steel sheet suitable for automotive outer panels. More particularly, the present invention relates to a high-strength cold-rolled steel sheet, which has an r₉₀ of 1.3 or less, an average plastic strain ratio r_(m) near 1, a low in-plane anisotropy index □r of 0.15 or less, thereby providing excellent shape-fixability so as to allow isotropic plastic deformation of the steel sheet during press forming, and a method for manufacturing the same.

BACKGROUND ART

As for steel plates for automobile bodies, cold-rolled steel sheets of excellent formability have been required to perform press forming without work defects and to efficiently manufacture automotive components of a desired shape after the press forming. In particular, for automotive outer panels, it is necessary to have dent resistance and shape-fixability.

In view of dent resistance, bake hardened steel sheets have been generally applied, the strength of which is increased after painting. In addition, in order to enhance the shape-fixability, the panels must be uniformly deformed in a plane direction thereof, and subjected to a lower load during the press forming.

For automotive inner panels, since the panels are generally subjected to deformation in a deep drawing mode, it is advantageous to provide cold-rolled steel sheets which have high elongation and plastic strain ratio during processing the panels.

On the other hand, for the automotive outer panels, since the panels are generally subjected to deformation in a stretching mode, it is advantageous to provide cold-rolled steel sheets which are subjected to uniform deformation in the plane direction and have low biaxial yield strength. With such cold-rolled steel sheets having excellent shape-fixability so as to permit uniform plastic deformation in the plane direction and low biaxial yield strength, the automotive outer panels which have a complicated shape can be advantageously produced.

Elongation is one of the mechanical properties of a material, and measures a percentage change in length of the material which elongates without fracture when tensile force is applied to the material. Thus, high elongation of a steel sheet permits large deformation of the steel sheet.

Plastic strain ratio “r” is a value which can be defined by a ratio of strain in a width direction to strain in a thickness direction. A high plastic strain ratio means that, assuming that a steel sheet has a constant strain amount in the width direction, a steel sheet with a high plastic strain ratio has a low strain in the thickness direction when applying tensile force to the steel sheet by a predetermined deformation amount in a certain direction, and thus the steel sheet can be worked without necking even with a large deformation amount. The plastic strain ratio is caused by anisotropic properties of the steel sheet, and thus exhibits different values according to tensile directions.

As for values measuring a degree of change in plastic strain ratio according to the tensile directions, an average plastic strain ratio r_(m) and an in-plane anisotropy index □r are provided, which can be calculated by the following Equations (1) and (2), respectively: r _(m)=(r ₀+2r ₄₅ +r ₉₀)/4  (1) □r=(r ₀−2r ₄₅ +r ₉₀)/2  (2)

wherein r₀, r₄₅, and r₉₀ are plastic strain ratios in tensile directions at 0, 45, and 90 degrees with respect to a rolling direction on the steel sheet, respectively.

A high plastic strain ratio results in high biaxial yield strength, and makes it difficult to work the outer panels of the complicated shape. FIG. 1 shows theoretical results based on Taylor polycrystal modeling as to influence of plastic strain ratio on a locus of the yield strength of steel which comprises two different major textures. In the case of Interstitial Free (IF) steel having a high plastic strain ratio, it can be seen from FIG. 1 that, even if the yield strength of the steel in a rolling direction is the same as that of isotropic steel having an average plastic strain ratio r_(m)=1, the IF steel has a high biaxial yield strength.

Thus, in order to reduce the biaxial yield strength, it is preferable that r_(m) be lowered to near 1.

In addition, in order to lower □r, it is necessary to reduce the difference between the plastic strain ratios according to the respective directions when tensile force is applied to the steel in the respective directions.

In other words, the low □r means that the distribution of strain is uniform in the plane direction of the steel sheet during the press forming, and is advantageous for forming of the steel sheet while leading to uniform deformation thereof in a stretching mode. As such, steel having an r_(m) approaching 1 and a low □r enhances the shape fixability during work for the automotive outer panels which will be subjected to major deformation in the stretching mode.

Other well-known techniques for enhancing the formability of the automotive steel sheet will be described hereinafter.

According to a technique for enhancing the formability of an automotive steel sheet disclosed in Japanese Patent Laid-open Publication No. (Hei) 9-296226, Ti or Nb is added as a single component or a mixture thereof to an ultra-low carbon cold-rolled steel sheet, and solid-solutions C and N are precipitated as a carbide and nitride to improve the elongation and the plastic strain ratio, thereby enhancing formability.

In addition, according to techniques for enhancing the formability disclosed in Japanese Patent Laid-open Publication Nos. (Hei) 6-158176, (Hei) 8-109416, (Hei) 11-40531, (Hei) 4-95392, and 2002-3951, the in-plane anisotropy of the steel sheet is reduced to lower defects such as plane defects during the press forming. According to the conventional techniques mentioned above, the in-plane anisotropy of the ultra-low carbon cold-rolled steel sheet is lowered through grain refinement of hot-rolled structures using a quenching apparatus immediately after finishing mill.

However, the conventional techniques have a problem in that, since the Ti and/or Nb-added ultra-low carbon steel has relatively high r_(m) and □r, it exhibits severe in-plane anisotropy and high biaxial yield strength for the deformation in the stretching mode, and thus is disadvantageous in terms of shape fixability even though it exhibits excellent formability in a deep drawing mode. In addition, according to the conventional techniques, since only 0.005% or less of carbon is generally added for enhancing deep drawability, it is difficult to obtain high strength.

Meanwhile, DE 3843732, DE 3803064, and U.S. Pat. No. 5,139,580 disclose a method for manufacturing high strength cold-rolled steel sheets having isotropic plasticity by controlling a carbide and fine textures during hot-rolling and annealing through addition of Ti, Nb, V, and the like, which are carbide formation elements in the low carbon steel.

However, the conventional techniques have problems in that these techniques are performed using a batch annealing apparatus, and require a long period of time for the process, thereby lowering productivity per unit time.

In addition, Japanese Patent Laid-open Publication No. (Hei) 10-130780 discloses a technique for manufacturing high-strength isotropic steel from Ti or Nb added low-carbon steel using a continuous annealing apparatus. The purpose of this technique is to manufacture steel sheets of a low □r using strong correlation between recrystallization total elongation and □r of the Ti or Nb-added steel sheet.

In other words, according to this conventional technique, although the steel sheets are formed to have □r of 0.1 or less in order to reduce ear formation when manufacturing automotive components formed to have circular or angular column shapes, the fact that excellent shape fixability can be ensured by r_(m) lowered to near 1 is not considered.

Meanwhile, Aging Index (AI) of 30 MPa or less is required for typical steel sheets for automotive outer panels. In this regard, since the conventional technique is applied to the steel sheet formed to have the circular or angular column shape, no consideration is paid to aging index.

Meanwhile, U.S. Pat. No. 6,162,308 discloses a technique for manufacturing a high-strength isotropic steel sheet from Ti and/or Nb added low-carbon steel using a continuous annealing apparatus. Since the purpose of the conventional technique is to manufacture a non-aging low-carbon steel sheet which does not require overaging, it is necessary to add at least one of Cu, V and Ni up to an amount of 0.15% in addition to Ti and Nb. In addition, since the conventional steel sheet has □r in the range of 0.15˜0.28, it is undesirable in view of isotropy.

DISCLOSURE OF INVENTION

Technical Problem

The present invention has been made to solve the above problems, and it is an object of the present invention to provide a high-strength isotropic cold-rolled steel sheet, which is made using a low-carbon steel comprising a little amounts of Ti to have an r₉₀ of 1.3 or less, an r_(m) approaching 1, and a □r of 0.15 or less, ensuring excellent shape-fixability so as to be suitable for steel for automotive outer panels which are essentially subjected to deformation in a stretching mode, and a method for manufacturing the same.

Technical Solution

In accordance with one aspect of the present invention, the above and other objects can be accomplished by the provision of a high-strength cold-rolled steel sheet having excellent shape-fixability, comprising: 0.01˜0.05% of C; 0.005˜0.06% of Ti; 0.1˜1% of Mn; 0.1% or less of Si; 0.03% or less of P; 0.03% or less of S; 0.08% or less of Sol. Al; 0.01% or less of N; the balance of Fe and other unavoidable impurities in terms of weight %, wherein a composition of Ti and N satisfies the relationship: Ti/N>5, and a composition of Ti and C satisfies the relationship: (48/12)C−Ti*>0.03% [where Ti*=Ti−(48/14)N], and the steel sheet has an Aging Index (AI) of 30 MPa or less.

More preferably, the steel sheet comprises 0.015˜0.035% of C.

More preferably, the steel sheet comprises 0.01˜0.04% of Ti.

More preferably, the relationship (48/12)C−Ti* is in the range of 0.06˜0.11%.

In accordance with another aspect of the present invention, a method for manufacturing a high-strength cold-rolled steel sheet having excellent shape-fixability, the method comprising the steps of: finish rolling steel at an Ar₃ temperature or more to provide a hot rolled steel sheet, the steel comprising: 0.01˜0.05% of C; 0.005˜0.06% of Ti; 0.1˜1% of Mn; 0.1% or less of Si; 0.03% or less of P; 0.03% or less of S; 0.08% or less of Sol. Al; 0.01% or less of N; the balance of Fe and other unavoidable impurities in terms of weight %, wherein a composition of Ti and N satisfies the relationship: Ti/N>5, and a composition of Ti and C satisfies the relationship: (48/12)C−Ti*>0.03% [where Ti*=Ti−(48/14)N], the steel having an aging index of 30 MPa or less; rapid quenching the hot rolled steel sheet at a rate of 50° C./sec or more, followed by coiling the steel sheet at a temperature of 650° C. or less; acid pickling the coiled steel sheet, followed by cold rolling the steel sheet at a reduction rate of 50˜80%; annealing the cold-rolled steel sheet after heating to a temperature of recrystallization temperature ˜Ac₃; primary cooling of the annealed steel sheet to a temperature of 600˜700° C. at a rate of 3° C./sec or more, followed by secondary cooling of the cooled steel sheet to a temperature of 100˜500° C. at a rate of 30° C./sec or more; overaging the cooled steel sheet at a temperature of 200˜500° C. within 10 minutes; and skin pass rolling the steel sheet at a reduction rate of 0.5% or more.

More preferably, the steel comprises 0.015˜0.035% of C.

More preferably, the steel comprises 0.01˜0.04% of Ti.

More preferably, the relationship (48/12)C−Ti* is in the range of 0.06˜0.11%.

More preferably, rapid quenching of the hot rolled steel sheet is performed within 1 second of completion of finish rolling.

More preferably, annealing of the steel sheet is performed at a temperature of 760˜820° C. for 5 minutes or less.

More preferably, the steel sheet is heated at a rate of 3° C./sec or more for annealing.

Advantageous Effects

As apparent from the above description, the present invention can provide a high-strength isotropic cold-rolled steel sheet, which has an r₉₀ of 1.3 or less, an r_(m) approaching 1, and a low □r of 0.15 or less, ensuring excellent shape-fixability so as to be suitable for steel for automotive outer panels which are essentially subjected to deformation in a stretching mode, and the method for manufacturing the same. With the steel sheet of the invention, it is possible to easily process complicated automotive components in the stretching mode when forming the automotive components.

BRIEF DESCRIPTION OF THE DRAWINGS

The above and other objects, features and other advantages of the present invention will be more clearly understood from the following detailed description taken in conjunction with the accompanying drawings:

FIG. 1 is a diagram illustrating a relationship between a plastic strain ratio and a locus of yield strength;

FIG. 2 is a diagram illustrating a continuous annealing process in accordance with one embodiment of the present invention, and change in microstructure by the continuous annealing process;

FIG. 3 is a diagram illustrating components of major textures developing in steel;

FIG. 4 is a diagram illustrating influence of the textures on an r value;

FIG. 5 is a crystallographic orientation map of Inventive steel A obtained using an EBSD apparatus after continuous annealing of the Inventive steel A;

FIG. 6 is an optical micrograph obtained after continuous annealing of the Inventive steel A; and

FIG. 7 shows a cross-section of the Inventive steel A obtained by an orientation distribution function of Ψ₂=45° measured after continuous annealing of the Inventive steel A.

BEST MODE FOR CARRYING OUT THE INVENTION

Preferred embodiments will now be described in detail with reference to the accompanying drawings.

The inventors of the present invention have theoretically found that as an r_(m) of steel is lowered to near 1, the biaxial yield strength of the steel is also lowered, thereby providing excellent shape-fixability to the steel. Then, various investigations were continuously carried out by the inventors in order to provide a technique for manufacturing a cold-rolled steel sheet using low-carbon steel comprising a little amounts of Ti to have excellent shape-fixability, isotropic structure, and an aging index of 30 MPa or less such that the steel sheet can be suitably used for automotive outer panels.

As a result, it has been found that, by adjusting the amounts of Ti, N and C in the steel sheet to have predetermined compositional relationships therebetween, and conditions for manufacturing the steel sheet, in particular, a hot rolling condition and an annealing condition, the high-strength cold-rolled steel sheet having an r₉₀ of 1.3 or less, an r_(m) approaching 1, a low □r of 0.15 or less, and an aging index of 30 MPa can be manufactured through a continuous annealing apparatus.

At first, the compositions of the steel sheet according to the invention, and the reasons for the range thereof will now be described in detail in which “weight %” will simply be represented as “%” hereinafter.

Carbon is an interstitial solid solution element in steel, and has a very significant influence on strength and texture of a steel sheet during cold rolling and annealing while existing in the form of cementite. The carbon content is preferably in the range of 0.01˜0.05%. When the carbon content is less than 0.01%, the steel sheet is lowered in strength, and excessively increased in □r.

Thus, it is desirable that the carbon content be 0.01% or more. In addition, since C is coupled with Fe to form the cementite in the steel, C can be stably present in the steel. According to the present invention, in order to avoid room temperature aging, it is necessary to have an appropriate amount of C such that C is precipitated to the cementite in the steel. Since an excessive amount of C causes a significant increase in strength, and reduction in ductility of the steel so that cold rolling properties of the steel is deteriorated, it is preferable that the upper limit of carbon content be 0.05% or less.

More preferably, the carbon content is in the range of 0.015˜0.035%. When the steel is heated in continuous annealing, C is coupled with Ti to precipitate TiC in the steel. The precipitated TiC provides precipitation hardening effect to the steel, which results in an increase in the strength of the steel. In addition, Normal Direction (ND) advantageous for reduction in □r servers to extend recovery and recrystallization rates of crystal grains which have a crystallographic orientation in parallel to <111> (<111>//ND), so that a fraction of the crystal grains having a crystallographic direction of <111>//ND is lowered. At this time, trace amounts of C are precipitated to Ti₄C₂S₂ at high temperature, which is coarser than TiC, and thus has substantially no influence on development in crystallographic orientation of recrystalline grains.

According to the present invention, Ti is one of the most important elements in addition to C. Ti is coupled with N as well as C to form TiN, and provides effects of suppressing formation of AlN. AlN precipitates formed during hot rolling cause elongation of a hot rolled structure, thereby increasing shape anisotropy of the steel sheet. As such, Ti serves to lower a fraction of crystal grains having an orientation of strong anisotropy by suppressing formation of AlN while precipitating TiC, and thus has effect of lowering □r and increasing the strength of the steel by virtue of precipitation hardening.

However, since Ti is an expensive element, it is advantageous in view of manufacturing costs to add as little Ti to the steel as possible. Thus, according to the present invention, preferably, Ti is in the range of 0.005˜0.06% under consideration of manufacturing costs without deteriorating the effects obtained by addition of Ti. More preferably, Ti is in the range of 0.01˜0.04%. At this time, in order to permit Ti to be precipitated to TiC during annealing while suppressing formation of AlN, Ti must be added to the steel such that a ratio of Ti to N (Ti/N) is more than 5, i.e. Ti/N>5.

In addition, since it is necessary to have an appropriate amount of C such that the appropriate amount of C is precipitated to the cementite in order to suppress the room temperature aging, it is preferable that Ti content satisfy the following relationship with respect to C and N: (48/12)C−Ti*>0.03%, where Ti*=Ti−(48/14)N.

Here, Ti* refers to an effective Ti content, which is an amount of Ti necessary to form TiC excluding an amount of Ti necessary to form TiN in order to suppress the formation of AlN during hot rolling. More preferably, a ratio of C to the effective Ti (Ti*), i.e. (48/12)C−Ti*, is in the range of 0.06˜0.11%.

Mn is an effective element for solid solution strengthening in steel, and precipitates S of the steel to MnS, thereby suppressing slip breakage and high temperature embrittlement caused by S during hot rolling. According to experimentation with respect to the invention, when Mn content is less than 0.1%, increased strength cannot be obtained, and S cannot be precipitated by Mn, thereby making it difficult to ensure the formability of the steel. On the contrary, when Mn content is more than 1%, advantageous effects caused by addition of Mn are saturated.

Si serves as a solid solution strengthening element in steel, and is preferably added to an amount of 0.1% or less in order to ensure proper elongation of the steel.

Higher P content is very advantageous for an increase in strength of steel. However, an excessive amount of P increases possibility of brittle fracture of the steel, which results in a high possibility of slip breakage of a slab during hot rolling. In addition, when the excessive amount of P is added to the steel, P is easily diffused into gain boundaries and segregated therein after annealing, thereby causing secondary work embrittlement during a forming process. Thus, it is necessary to restrict the P content. According to the present invention, since a desired strength of the steel can be obtained by the precipitation hardening effect caused by TiC, the P content is preferably restricted to 0.03% or less.

Nitrogen and sulfur are unavoidable elements introduced into steel during a steel manufacturing process, and thus it is important to keep contents of N and S as low as possible. However, in order to ultimately lower the contents of N and S, there is a problem of increasing refining costs of the steel. In this regard, it is desirable that the contents of N and S be lowered within an acceptable range of operating conditions, and according to the invention, S content is preferably restricted to 0.03% or less. In addition, since N forms TiN at a high temperature, and thus changes an effective content of Ti which will be coupled with C, higher N content causes a problem of lowering the effective Ti content. Thus, according to the invention, N content is preferably restricted to 0.01% or less.

Sol. Al effectively serves as a deoxidation element of molten steel. However, since an excessive amount of Sol. Al can have a negative influence on the formability of the steel, the content of Sol. Al is preferably restricted to 0.08% or less.

A method for manufacturing a high strength steel sheet according to the present invention will be described in detail as follows.

As a raw material of a hot rolled steel sheet, steel formed to have the above composition through continuous casting may be used without being formed into an ingot. Alternatively, a steel ingot having the above composition may be used after being reheated. At this time, when forming the steel sheet using the ingot, it is desirable that the ingot be reheated to 1,200° C. or more such that Ti₄C₂S₂ formed during cooling of the ingot is resolved in the steel as a solid solution.

According to the invention, hot rolling is performed to form a hot rolled steel sheet according to a typical process, and it is desirable that a final pass of finish rolling be terminated at a temperature of Ar₃ or more. If the final temperature of hot rolling is lowered, the surface and the edges of the hot rolled steel sheet are hot-rolled at a temperature of two-phase region so that crystal grains become coarse and non-uniform, causing surface defects of the steel sheet during press forming.

After finish rolling, the steel sheet is rapid quenched at a rate of 50° C./sec or more to a coiling temperature or more on a Run Out Table (ROT) so as to form fine crystal grains in the hot rolled steel sheet. If the steel sheet is quenched at a rate less than 50° C./sec, crystal grains become coarse.

In addition, more preferably, the steel sheet is quenched within 1 second of completion of finish rolling so as to form finer crystal grains. Rapid quenching of the steel sheet can be performed using a high density cooler equipped in front of the ROT. After quenching the steel sheet, the steel sheet is preferably coiled at a temperature of 650° C. or less. The reason being that a coiling temperature exceeding 650° C. causes coarsening of TiC precipitates, which weakens the function of delaying recovery and recrystallization rate of sub-grains which have an orientation of strong anisotropy during annealing, thereby increasing a fraction of crystal grains having the orientation of strong anisotropy.

The coiled steel sheet is acid-pickled by a typical process, and is then preferably subjected to cold rolling at a reduction rate of 50˜80%. If the reduction rate of cold rolling is less than 50%, recrystallization does not sufficiently occurred during annealing, thereby lowering ductility, and if the reduction rate of cold rolling is more than 80%, the in-plane anisotropy of the steel sheet is increased.

According to the invention, annealing refers to continuous annealing as shown in FIG. 2, and is performed at a recrystallization temperature or more, and at a temperature less than Ac₃ or less. If the annealing temperature exceeds Ac₃, the steel sheet is annealed in a two-phase coexistence region of α and γ, so that recrystallization and grain growth of the crystal grains having the orientation of strong anisotropy are promoted, causing coarsening of the crystal grains. Since coarsened crystal grains cause deterioration in strength and ductility at the same time, the annealing temperature is preferably restricted to a temperature of Ac₃ or less.

On the other hand, if the annealing temperature is significantly lowered below the recrystallization temperature, ductility is deteriorated. Thus, more preferably, the annealing temperature is in the range of 760˜820° C. In addition, it is preferable that the steel sheet be maintained within 5 minutes at the annealing temperature described above. The reason being that longer maintenance of the steel sheet at the annealing temperature described above causes an increase in r₉₀, and growth of crystal grains having the crystallographic orientation of strong anisotropy.

In addition, the cold-rolled steel sheet is preferably heated to the annealing temperature at a rate of 3° C./sec or more. The reason being that a heating rate less than 3° C./sec causes an increase in annealing period, thereby possibly coarsening the crystal grains.

Thereafter, primary cooling of the annealed steel sheet is performed to a temperature of 600˜700° C., which is a temperature providing high solid solubility of C in an Fe matrix, and then secondary cooling is immediately performed to a temperature of 100˜500° C., which is a temperature of providing low solid solubility of C in the Fe matrix, so as to lead precipitation of cementite in grain boundaries and interfaces. Primary cooling is preferably performed at a rate of 3° C./sec or more. Meanwhile, secondary cooling is preferably performed at a rate of 30° C./sec or more.

The reason being that, if the secondary cooling rate is less than 30° C./sec, oversaturated C prevents the cementite from being sufficiently precipitated, thereby causing deterioration in ductility and room temperature aging. After secondary cooling, the steel sheet is reheated to a temperature in the range of 200˜500° C., followed by overaging for 10 minutes or less so as to permit growth of the precipitated cementite. If overaging is performed at a temperature less than 200° C., the cementite is not sufficiently grown so that a portion of C is resolved as the solid solution, deteriorating the ductility and room temperature aging, and if averaging is performed at a temperature more than 500° C., the solubility of C in the Fe matrix is increased, thereby deteriorating the ductility and room temperature aging.

Then, skin pass rolling is preferably performed upon the steel sheet at a reduction rate of 0.5% or more.

MODE FOR THE INVENTION

The following examples are provided for a more clear understanding of the present invention. It should be understood that specific processes, conditions, ratios, and reported data for illustrating principle and examples of the present invention are exemplary, and thus do not limit the scope of the present invention.

After melting and continuous casting Ti-added low carbon steel having compositions as shown in Table 1, the steel was reheated to a temperature of 1,200° C., and subjected to finish hot rolling at a temperature of 870˜890° C. to 2.5 □. Then, after passing a quenching start time shown in Table 2, the hot rolled steel was quenched at a rate of 60° C./sec through a high density cooler, and coiled at a coiling temperature shown in Table 2. After removing surface oxidation layer from the coiled steel sheet, the steel sheet was cold-rolled at a reduction rate of 70% to 0.75 □.

Then, the cold-rolled steel sheet was subjected to heat treatment on a continuous annealing line. During the heat treatment, the steel sheet was heated to a maximum temperature of 780˜800° C. After heating the steel sheet at this temperature for 1 minute, the steel sheet was primarily cooled to 700° C. at a rate of 5° C./sec, and then secondarily cooled to 100° C. at a rate of 60° C./sec. The steel sheet was reheated to a temperature of 300˜350° C., and was subjected to overaging for 3 minutes, and skin pass rolling at a reduction rate of 1˜1.3%. Tensile test of the annealed sheet obtained by the above processes was performed after processing the annealed sheet to EN10002-1 test samples. TABLE 1 Steel Composition (wt %) Steel C Si Mn P S Sol-A1 Ti N(ppm) Ti/N Ti*(ppm) 4C-Ti* A 0.019 0.012 0.19 0.011 0.011 0.023 0.019 32 5.94 81 0.068 B 0.025 0.012 0.20 0.011 0.012 0.041 0.021 30 7.00 107 0.110 C 0.028 0.011 0.19 0.018 0.011 0.046 0.015 27 5.56 57 0.106 D 0.019 0.012 0.19 0.011 0.011 0.023 0.019 32 5.94 80 0.068 E 0.02 0.011 0.19 0.014 0.011 0.039 0.018 26 6.92 91 0.071 F 0.024 0.012 0.20 0.011 0.012 0.041 0.021 30 7.00 107 0.085 G 0.021 0.006 0.20 0.011 0.012 0.044 0.022 26 8.46 131 0.071 H 0.020 0.008 0.19 0.011 0.012 0.035 0.029 26 11.15 201 0.060 I 0.018 0.011 0.19 0.011 0.008 0.027 0.013 31 4.19 23 0.068 J 0.024 0.079 0.19 0.010 0.012 0.035 0.015 31 4.84 44 0.092 K 0.019 0.013 0.19 0.011 0.012 0.03 0.014 32 4.38 30 0.073 L 0.023 0.049 0.19 0.010 0.011 0.033 0.015 32 4.69 40 0.088 M 0.016 0.014 0.2 0.011 0.012 0.045 0.021 29 7.24 111 0.053 N 0.021 0.016 0.19 0.011 0.012 0.042 0.040 27 14.8 307 0.053 O 0.028 0.014 0.2 0.011 0.012 0.037 0.021 31 6.77 105 0.102

The following Table 2 shows conditions of manufacturing a cold-rolled steel sheet having the composition shown in Table 1, and results of a uni-axial test. In Table 2, FDT indicates final temperature of finish rolling, CT indicates a coiling temperature, ST indicates an annealing temperature, YP indicates a yield strength, TS indicates a tensile strength, El indicates total elongation, r₉₀ indicates a plastic strain ratio in a directions of 90 degrees with respect to a rolling direction of the steel sheet, □r indicates an in-plane anisotropy index, and AI indicates an aging index. AI was calculated using a difference between flow stress after application of 7.5% pre-strain before heating the steel sheet and flow stress after heating the steel sheet at 100° C. for 1 hour. TABLE 2 Manufacturing conditions Quenching Mechanical properties FDT start CT ST YP TS AI Steel (° C.) time(sec) (° C.) (° C.) (Mpa) (Mpa) E1(%) r90 |Δr| (Mpa) IS A 892 0.3 500 780 228 334 39.2 1.21 0.14 26 B 890 0.3 600 780 220 340 36.2 1.09 0.1 27 C 883 0.3 500 790 254 350 36.8 1.25 0.13 26 D 891 0.3 500 780 233 335 40.2 1.14 0.11 25 E 893 0.3 500 780 252 354 37.5 1.27 0.12 23 F 892 0.3 600 780 220 340 37 0.97 0.07 27 G 890 0.3 540 800 284 360 33.7 1.07 0.08 24 H 883 0.3 540 800 279 355 33.6 1.04 0.1 28 CS I 887 1.3 600 780 218 324 40.3 1.29 0.26 35 J 889 1.3 500 780 242 339 38.4 1.27 0.26 33 K 888 0.3 600 780 220 327 41.2 1.22 0.26 22 L 886 0.3 500 780 220 340 37.2 1.22 0.21 23 M 882 0.3 700 780 235 335 37.4 1.31 0.21 34 N 886 1.3 540 800 285 376 31.8 0.99 0.01 39 O 885 1.3 500 800 272 356 34.64 1.18 0.18 33

In Table 2, Inventive steels A˜H satisfy the composition and manufacturing conditions of the present invention. As can be seen from Table 2, these samples have an r₉₀ of 1.3 or less, and a □r of 0.15 or less, thereby providing a low biaxial yield strength and a low in-plane anisotropy index.

Meanwhile, Comparative steels I˜L deviate from the range of the present invention, and have a low Ti content relative to an N content. In other words, since Ti/N is lower than 5 which is in the range of the present invention, these comparative steels has a □r of 0.15 or more. In particular, Comparative steels I and J have manufacturing conditions wherein quenching start time thereof after finish rolling is longer than that of the present invention.

In Comparative steels M and N, Ti/N ratios are in the range of the present invention, whilst r₉₀, □r, and aging index thereof are not in the range of the invention. For Comparative steel M, it is considered that, since a coiling temperature is higher than that of the invention, allowing TiC to be precipitated by solid solution C and coarsened in the hot rolled sheet, precipitation of TiC is insufficient during annealing, so that development of crystallographic orientation ({554}<225>) having a high r₉₀ and □r is increased, thereby failing to obtain isotropic steel. For Comparative steels M and N, it is considered that since Ti content does not satisfy the relationship (48/12)C−Ti*=0.6, the steels has a high aging index.

In addition, for Comparative steel O, a quenching start time is not in the range of the present invention. In comparison to a shortened quenching start time, the structure of the hot rolled steel sheet is coarsened, reducing the number of nucleation sites for cementite during cooling after annealing. Thus, Comparative steel O has a high room temperature aging index, and a □r of 0.15 or more.

FIG. 3 shows an orientation distribution function of Ψ₂=45° with respect to major components of texture developing in the steel. FIG. 4 shows results of theoretical calculation for influence of texture on anisotropy in plastic strain ratio of major components of the texture shown in FIG. 3 using Taylor polycrystal theory.

From FIG. 4, it can be appreciated that texture of α-fibre (RD//<110>) and texture of γ-fibre (ND//<111>) have different influences on the plastic strain ratio. For the texture of α-fibre (RD//<110>), the plastic strain ratio is generally low, and r₄₅ is the highest value, whereas for texture of {554}<225>including γ-fibre (ND//<111>), r₄₅ is the lowest value. As such, it can be appreciated that suitable combination of texture described above is required to provide isotropic steel.

FIG. 5 is a Crystallographic Orientation Map (COM) of Inventive steel A obtained using an Electron Back-Scattered Diffraction (EBSD) apparatus attached to a Field Emission Scanning Electron Microscope (FE-SEM). When comparing colors on the inverse pole figure of FIG. 5, it can be seen that both textures of α-fibre (RD//<110>) and γ-fibre (ND//<111>) are sufficiently developed for the inventive steel. FIG. 6 shows results of analyzing crystal grains and cementite via optical microscopy. It can be seen from FIG. 6 that the cementite is mainly formed in the grain boundaries. FIG. 7 shows an Orientation Distribution Function (ODF) of Ψ₂=45° based on data obtained by correcting data of a pole figure which is provided by X-ray diffraction for micro-texture developing in the Inventive steel A.

It can be seen from FIG. 7 that both textures of α-fibre (RD//<110>) and γ-fibre (ND//<111>) are sufficiently developed. According to the test results, it can be appreciated that, since both textures of α-fibre (RD//<110>) and γ-fibre (ND//<111>) are sufficiently developed in the inventive steels, the inventive steels have excellent isotropy.

It should be understood that the embodiments and the accompanying drawings have been described for illustrative purposes, and the present invention is limited only by the following claims. Further, those skilled in the art will appreciate that various modifications, additions and substitutions are allowed without departing from the scope and spirit of the invention according to the accompanying claims. 

1. A high-strength cold-rolled steel sheet having excellent shape-fixability, comprising in percent by weight: 0.01˜0.05% of C; 0.005˜0.06% of Ti; 0.1˜1% of Mn; 0.1% or less of Si; 0.03% or less of P; 0.03% or less of S; 0.08% or less of Sol. Al; 0.01% or less of N; and the balance of Fe and other unavoidable impurities in terms of weight %, wherein a composition of Ti and N satisfies the relationship: Ti/N>5, and a composition of Ti and C satisfies the relationship: (48/12)C−Ti*>0.03% [where Ti*=Ti−(48/14)N], the steel sheet having an Aging Index (AI) of 30 MPa or less.
 2. The cold-rolled steel sheet according to claim 1, wherein the steel sheet comprises 0.015˜0.035% of C.
 3. The cold-rolled steel sheet according to claim 1, wherein the steel sheet comprises 0.01˜0.04% of Ti.
 4. The cold-rolled steel sheet according to claim 1, wherein the relationship (48/12)C−Ti* is in the range of 0.06˜0.11%.
 5. A method for manufacturing a high-strength cold-rolled steel sheet having excellent shape-fixability, the method comprising the steps of: finish rolling steel at an Ar₃ temperature or more to provide a hot rolled steel sheet, the steel comprising in percent by weight: 0.01˜0.05% of C; 0.005˜0.06% of Ti; 0.1˜1% of Mn; 0.1% or less of Si; 0.03% or less of P; 0.03% or less of S; 0.08% or less of Sol. Al; 0.01% or less of N; and the balance of Fe and other unavoidable impurities in terms of weight %, wherein a composition of Ti and N satisfies the relationship: Ti/N<5, and a composition of Ti and C satisfies the relationship: (48/12)C−Ti*<0.03% [where Ti*=Ti−(48/14)N], the steel having an aging index of 30 MPa or less; rapid quenching the hot rolled steel sheet at a rate of 50° C./sec or more, followed by coiling the steel sheet at a temperature of 650° C. or less; acid pickling the coiled steel sheet, followed by cold rolling the steel sheet at a reduction rate of 50˜80%; annealing the cold-rolled steel sheet after heating to a temperature of recrystallization temperature˜AC₃; primary cooling of the annealed steel sheet to a temperature of 600˜700° C. at a rate of 3° C./sec or more, followed by secondary cooling of the cooled steel sheet to a temperature of 100˜500° C. at a rate of 30° C./sec or more; overaging the cooled steel sheet at a temperature of 200˜500° C. within 10 minutes; and skin pass rolling the steel sheet at a reduction rate of 0.5% or more.
 6. The method according to claim 5, wherein the steel comprises 0.015˜0.035% of C.
 7. The method according to claim 5, wherein the steel comprises 0.01˜0.04% of Ti.
 8. The method according to claim 5, wherein the relationship (48/12)C−Ti* is the range of 0.06˜0.11%.
 9. The method according to claim 5, wherein rapid quenching of the hot rolled steel sheet is performed within 1 second of completion of finish rolling.
 10. The method according to claim 5, wherein annealing of the steel sheet is performed at a temperature of 760˜820° C. for 5 minutes or less.
 11. The method according to claim 5, wherein the steel sheet is heated at a rate of 3° C./sec or more for annealing.
 12. The cold-rolled steel sheet according to claim 2, wherein the relationship (48/12)C−Ti* is in the range of 0.06˜0.11%.
 13. The cold-rolled steel sheet according to claim 3, wherein the relationship (48/12)C−Ti* is in the range of 0.06˜0.11%.
 14. The method according to claim 6, wherein the relationship (48/12)C−Ti* is the range of 0.06˜0.11%.
 15. The method according to claim 7, wherein the relationship (48/12)C−Ti* is the range of 0.06˜0.11%.
 16. The method according to claim 10, wherein the steel sheet is heated at a rate of 3° C./sec or more for annealing. 